JP6935579B2 - Cobalt-based alloy product and method for manufacturing the product - Google Patents

Description

The present invention relates to a cobalt-based alloy material having excellent mechanical properties, and more particularly to a cobalt-based alloy product, a method for producing the product, and a cobalt-based alloy article.

Cobalt (Co) -based alloy material is a typical heat-resistant alloy material together with nickel (Ni) -based alloy material, and is also called a superalloy. Widely used for turbine members). Co-based alloy materials have been used as turbine stationary blades and turbine combustor members because they have higher material cost than Ni-based alloy materials, but have excellent corrosion resistance and abrasion resistance, and are easily solid-solved and strengthened. ..

Due to various alloy composition improvements and manufacturing process improvements that have been carried out to date in heat-resistant alloy materials, Ni-based alloy materials are _{strengthened by precipitation of the γ'phase (for example, Ni 3} (Al, Ti) phase). It was developed and is now mainstream. On the other hand, in the Co-based alloy material, since it is difficult to precipitate an intermetallic compound phase such as the γ'phase of the Ni-based alloy material, which greatly contributes to the improvement of mechanical properties, precipitation strengthening by the carbide phase has been studied.

For example, in Patent Document 1 (Japanese Patent Laid-Open No. 61-243143), massive and granular carbides having a particle size of 0.5 to 10 μm are precipitated on a base of a cobalt-based alloy having a crystal particle size of 10 μm or less. A Co-based superplastic alloy characterized by the above is disclosed. The cobalt-based alloy has a weight ratio of C: 0.15 to 1%, Cr: 15 to 40%, W and / or Mo: 3 to 15%, B: 1% or less, Ni: 0 to 20%, Nb: It is disclosed that it consists of 0 to 1.0%, Zr: 0 to 1.0%, Ta: 0 to 1.0%, Ti: 0 to 3%, Al: 0 to 3%, and the balance Co. According to Patent Document 1, Co-based superplasticity that exhibits superplasticity even in a low temperature region (for example, 950 ° C.), has an elongation rate of 70% or more, and can produce a complicated shape by plastic working such as forging. It is said that it can provide alloys.

According to Patent Document 2 (Japanese Patent Laid-Open No. 7-179967), in% by weight, Cr: 21 to 29%, Mo: 15 to 24%, B: 0.5 to 2%, Si: 0.1% or more and less than 0.5%, C : A Co-based alloy having excellent corrosion resistance, abrasion resistance, and high-temperature strength, which is more than 1% and 2% or less, Fe: 2% or less, Ni: 2% or less, and the balance substantially Co, is disclosed. According to Patent Document 2, the Co-based alloy has a composite structure in which molybdenum boride and chromium carbide are relatively finely dispersed in a quaternary alloy phase of Co, Cr, Mo, and Si, and has good corrosion resistance and resistance. It is said to have abrasion resistance and high strength.

By the way, in recent years, three-dimensional modeling technology (so-called 3D printing) such as additive manufacturing (AM method) has attracted attention as a technology for manufacturing a final product having a complicated shape by a near net shape, and the three-dimensional modeling technology has been attracting attention. Is being actively researched and developed to apply the above to heat-resistant alloy members.

For example, Patent Document 3 (Japanese Patent Laid-Open No. 2016-535169) provides a layer forming method including the following steps: a) a raw material for a powdery or suspension-like granular composite material having a void ratio of less than 20%. Steps to b) deposit the first part of the composite on the surface of the target, c) supply energy to the composite in the first part to sinter, fuse or blend the composite in the first part. The step of melting to form the first layer, d) the step of depositing the second portion of the composite material on the first layer, e) supplying energy to the composite material of the second part, said Disclosed is a step of sintering, fusing or melting the composite material in a second portion to form a second layer, wherein the energy is supplied by a laser.

According to Patent Document 3 (Special Table 2016-535169), selective laser melting (SLM) or direct metal laser melting (DMLM) generally involves one type of material (pure titanium or Ti-6Al-4V). It is said to be beneficial for such single alloys). In contrast, for multiple materials, alloys, or combinations of alloys with other materials such as plastics, ceramics, polymers, carbides, and glass, selective laser sintering (SLS) or direct metal laser sintering (DMLS). The law) is generally applied. Sintering is a technical concept different from melting, and the sintering process is a process in which the material powder is not completely melted, but is heated to the point where the powder can be fused together at the molecular level. That is.

Japanese Unexamined Patent Publication No. 61-243143 Japanese Unexamined Patent Publication No. 7-179967 Special Table 2016-535169

Since alloy members can be manufactured by 3D printing directly even if they have a complicated shape such as turbine blades, the manufacturing work time can be shortened and the manufacturing yield can be improved (that is, the manufacturing cost can be reduced). It is a very attractive technology from the viewpoint).

Co-based alloy materials as described in Patent Documents 1 and 2 are considered to have higher mechanical properties than the previous Co-based alloy materials, but are compared with recent precipitation-strengthened Ni-based alloy materials. Unfortunately, it cannot be said that it has sufficient mechanical properties. Therefore, at present, most of the research on laminated shaped bodies (AM bodies) for high temperature member applications is aimed at precipitation-strengthened Ni-based alloy materials.

However, in the AM body of the precipitation-strengthened Ni-based alloy, problems such as inhibition of the formation of the γ'phase, which is the key to mechanical properties, and internal defects in the product are likely to occur, which is expected as a result. There is a problem that sufficient mechanical properties cannot be obtained. This is because the current precipitation-strengthened Ni-based alloy materials used as high-temperature members are optimized on the premise of melting and casting processes in high vacuum, so the stage of preparing alloy powder for the AM method and AM It is considered that oxidation and nitriding of the Al component and Ti component constituting the γ'phase are likely to occur at the legal stage.

On the other hand, the Co-based alloy material as described in Patent Documents 1 and 2 does not presuppose the precipitation of an intermetallic compound phase such as the γ'phase of the Ni-based alloy material, and therefore easily oxidizes Al and Ti. It does not contain much and can be used for melting and casting processes in the air. Therefore, it is considered to be advantageous for the production of alloy powders for the AM method and the production of AM bodies. Further, the Co-based alloy material has an advantage of having corrosion resistance and wear resistance equal to or higher than that of the Ni-based alloy material.

However, as described above, the conventional Co-based alloy material has a weakness that the mechanical properties are lower than those of the γ'phase precipitation strengthened Ni-based alloy material. In other words, the mechanical properties equal to or higher than those of the γ'phase precipitation strengthened Ni-based alloy material (for example, creep rupture time of 1100 hours or more when the creep test is performed under the conditions of temperature 900 ° C. and stress 98 MPa) are achieved. If possible, the Co-based alloy AM body can be a very attractive high temperature member.

The present invention has been made in view of the above problems, and an object of the present invention is to provide a Co-based alloy product having mechanical properties equal to or higher than those of a precipitation-strengthened Ni-based alloy material and a method for producing the same. It is in. Another object of the present invention is to provide an article which is a base of a Co-based alloy product.

(I) One aspect of the present invention is a product made of a Co-based alloy.
The Co-based alloy is
With carbon (C) of 0.08% by mass or more and 0.25% by mass or less,
Boron (B) of 0.1% by mass or less and
Contains 10% by mass or more and 30% by mass or less of chromium (Cr)
It contains 5% by mass or less of iron (Fe) and 30% by mass or less of nickel (Ni), and the total of Fe and Ni is 30% by mass or less.
It contains tungsten (W) and / or molybdenum (Mo), and the total of the W and the Mo is 5% by mass or more and 12% by mass or less.
Contains one or more of titanium (Ti), zirconium (Zr), hafnium (Hf), vanadium (V), niobium (Nb) and tantalum (Ta) in a total of 0.5% by mass or more and 2% by mass or less.
With 0.5% by mass or less of silicon (Si),
With 0.5% by mass or less of manganese (Mn),
Containing with nitrogen (N) of 0.003% by mass or more and 0.04% by mass or less,
The rest consists of Co and impurities,
The impurities are
Aluminum (Al) of 0.5% by mass or less and
It has a chemical composition containing 0.04% by mass or less of oxygen (O), and has a chemical composition.
The product is a polycrystal of a parent phase crystal, and particles of the MC type carbide phase and particles of the M ₂₃ C ₆ type carbide phase are co-precipitated.
The particles of the MC-type carbide phase are dispersed and precipitated in the particles of the matrix crystal at an average interparticle distance of 0.13 μm or more and 2 μm or less.
The particles of the M ₂₃ C ₆ type carbide phase are precipitated on the grain boundaries of the matrix crystal.
It provides a Co-based alloy product characterized by the above.

The present invention can make the following improvements and modifications to the above-mentioned Co-based alloy product (I).
(I) The particles of the MC-type carbide phase are particles of the MC-type carbide phase containing the Ti, the Zr, the Hf, the V, the Nb and / or the Ta.
Particles of the M ₂₃ C ₆ type carbide phase, the Cr, the Fe, the W, particles of the Mo and / or M ₂₃ C ₆ type carbide phase containing the Mn.
(Ii) The chemical composition of the Co-based alloy is
When the Ti is contained, the Ti is 0.01% by mass or more and 1% by mass or less.
When the Zr is contained, the Zr is 0.05% by mass or more and 1.5% by mass or less.
When the Hf is contained, the Hf is 0.01% by mass or more and 0.5% by mass or less.
When the V is included, the V is 0.01% by mass or more and 0.5% by mass or less.
When the Nb is contained, the Nb is 0.02% by mass or more and 1% by mass or less.
When the Ta is contained, the Ta is 0.05% by mass or more and 1.5% by mass or less.
(Iii) The product has a creep rupture time of 1200 hours or more when a creep test is performed under the conditions of a temperature of 850 ° C. and a stress of 168 MPa.
(Iv) The product is a high temperature member.
(V) The high temperature member is a turbine blade, a turbine blade, a turbine combustor nozzle or a heat exchanger.

(II) Another aspect of the present invention is the above-mentioned method for producing a Co-based alloy product.
An alloy powder preparation step for preparing a Co-based alloy powder having the above chemical composition, and
An alloy powder bed preparation step of spreading the Co-based alloy powder to prepare an alloy powder bed having a predetermined thickness, and irradiating a predetermined region of the alloy powder bed with laser light to prepare the Co-based alloy powder in the region. A laser melt coagulation element step for local melting and rapid solidification, and a selective laser melting (SLM) step for forming a laminated model by repeating the steps.
The first heat treatment step of performing the first heat treatment in the temperature range of 750 ° C. or higher and lower than 1100 ° C. on the laminated model, and the first heat treatment step.
It has a second heat treatment step of subjecting the laminated model body subjected to the first heat treatment to a second heat treatment having a temperature lower than that of the first heat treatment in a temperature range of 600 ° C. or higher and 1000 ° C. or lower.
In the SLM step, the relationship between the predetermined thickness h (unit: μm) of the alloy powder bed, the output P (unit: W) of the laser beam, and the scanning speed S (unit: mm / s) of the laser beam. The predetermined thickness h, the output P, and the scanning speed S are set so that “15 <h <150” and “67 (P / S) −3.5 <h <2222 (P / S) +13” are satisfied. Control,
A method for producing a Co-based alloy product.

INDUSTRIAL APPLICABILITY The present invention can make the following improvements and changes in the above-mentioned method (II) for producing a Co-based alloy product.
(Vi) The alloy powder preparation step includes an alloy powder classifying step of classifying the Co-based alloy powder into a particle size range of 5 μm or more and 100 μm or less.

According to the present invention, it is possible to provide a Co-based alloy product having mechanical properties equal to or higher than those of a precipitation-strengthened Ni-based alloy material and a method for producing the same. Further, it is possible to provide an article which is a base of a Co-based alloy product.

It is a flow chart which shows the process example of the manufacturing method of the Co-based alloy product which concerns on this invention. It is a scanning electron microscope (SEM) observation image which shows an example of the microstructure of the Co-based alloy AM body obtained by a selective laser melting process. It is an SEM observation image which shows an example of the fine structure of the Co-based alloy laminated model which performed the 1st heat treatment. It is an SEM observation image which shows an example of the fine structure of the Co-based alloy product obtained in the 2nd heat treatment step. It is an example of the Co-based alloy product which concerns on this invention, and is the perspective schematic diagram which shows the turbine vane as a high temperature member. It is sectional drawing which shows an example of the gas turbine equipped with the Co-based alloy product which concerns on this invention. It is an example of the Co-based alloy product which concerns on this invention, and is the perspective schematic diagram which shows the heat exchanger as a high temperature member. It is an example of SLM conditions in a selective laser melting process, and is a graph showing the relationship between the thickness of an alloy powder bed and the amount of local heat input.

[Basic idea of the present invention]
As described above, in the Co-based alloy material, various research and development have been carried out for strengthening by precipitation of the carbide phase. Examples of the carbide phase that contributes to precipitation strengthening include MC-type carbide phases of Ti, Zr, Hf, V, Nb, and Ta, and composite carbide phases of these metal elements.

The Ti, Zr, Hf, V, Nb, and Ta components and the C component, which is essential for forming the carbide phase, are the final solidified parts (for example, dendrite boundaries and grain boundaries) during melt solidification of Co-based alloys. ) Has the property of being easily segregated. Therefore, in the conventional Co-based alloy material, the carbide phase particles are precipitated along the dendrite boundary of the matrix phase and the grain boundaries. For example, in ordinary castings of Co-based alloys, the average spacing and average crystal grain size of dendrite boundaries are usually on the ^{order of 10 1} to 10 ² μm, so the average spacing of carbide phase particles is also on the order of ^{10 1} to 10 ^{2 μm.} Become. Further, even in a process such as laser welding in which the solidification rate is relatively high, the average spacing of the carbide phase particles in the solidification portion is about 5 μm.

It is generally known that precipitation strengthening in alloys is inversely proportional to the average spacing between precipitates, and it is said that precipitation strengthening is effective when the average spacing between precipitates is about 2 μm or less. There is. However, in the above-mentioned conventional technique, the average spacing between the precipitates does not reach that level, and a sufficient effect of precipitation strengthening cannot be obtained. In other words, in the prior art, it was difficult to finely disperse and precipitate carbide phase particles that contribute to alloy strengthening. This is the main reason why Co-based alloy materials have been said to have insufficient mechanical properties compared to precipitation-strengthened Ni-based alloy materials.

As another carbide phase that can be precipitated in the Co-based alloy, there is a Cr carbide phase. Since the Cr component has high solid solubility in the Co-based alloy matrix and is difficult to segregate, the Cr carbide phase can be dispersed and precipitated in the matrix crystal grains. However, it is known that the Cr carbide phase has low lattice consistency with the Co-based alloy matrix crystal and is not so effective as a precipitation strengthening phase.

The present inventors can dramatically improve the mechanical properties of the Co-based alloy material if the carbide phase particles that contribute to precipitation strengthening can be dispersed and precipitated in the parent phase crystal grains in the Co-based alloy material. I thought I could do it. In addition, it was considered that a heat-resistant alloy material superior to the precipitation-strengthened Ni-based alloy material could be provided in combination with the good corrosion resistance and abrasion resistance originally possessed by the Co-based alloy material.

Therefore, the present inventors have diligently studied the alloy composition and the manufacturing method for obtaining such a Co-based alloy material. As a result, the Co group is optimized by optimizing the alloy composition and controlling the amount of heat input for local melting and rapid solidification within a predetermined range in the production using the AM method (particularly, the selective laser melting method). It has been found that minute-sized segregated cells in which specific components (components forming a carbide phase that contributes to alloy strengthening) are segregated are formed in the matrix crystal grains of the alloy material (AM body). Further, when the obtained AM body is subjected to a predetermined heat treatment, the particles of the MC-type carbide phase are finely dispersed and precipitated in the grains of the matrix crystal, and the M ₂₃ C _6- type carbide phase is formed on the grain boundaries of the matrix crystal. It has been found that a fine structure in which particles are precipitated can be obtained. The present invention has been completed based on this finding.

Hereinafter, embodiments according to the present invention will be described with reference to the drawings along with manufacturing procedures. However, the present invention is not limited to the embodiments taken up here, and can be appropriately combined with a known technique or improved based on the known technique without departing from the technical idea of the invention. ..

[Manufacturing method of Co-based alloy products]
FIG. 1 is a flow chart showing a process example of a method for producing a Co-based alloy product according to the present invention. As shown in FIG. 1, the method for producing a Co-based alloy product according to the present invention generally uses an alloy powder preparation step (S1) for preparing a Co-based alloy powder and the prepared Co-based alloy powder. A selective laser melting step (S2) for forming an AM alloy having a desired shape, a first heat treatment step (S3) for subjecting the formed AM alloy to a first heat treatment, and an AM alloy having undergone the first heat treatment. It has a second heat treatment step (S4) of performing a second heat treatment.

Further, although not shown in FIG. 1, the Co-based alloy product obtained in the second heat treatment step S4 is further subjected to a step of forming a corrosion-resistant coating layer and a step of surface finishing, if necessary. You may. The Co-based alloy powder obtained in the alloy powder preparation step S1 can be a Co-based alloy article according to the present invention.

Hereinafter, each step will be described in more detail.

(Alloy powder preparation process)
This step S1 is a step of preparing a Co-based alloy powder having a predetermined chemical composition. The chemical composition contains C of 0.08% by mass or more and 0.25% by mass or less, B of 0.1% by mass or less, and Cr of 10% by mass or more and 30% by mass or less, Fe is 5% by mass or less and Ni is 30% by mass. % Or less, Fe and Ni total of 30% by mass or less, W and / or Mo included, W and Mo total of 5% by mass or more and 12% by mass or less, Ti, Zr, Hf, V , Nb and Ta of 1 or more in total of 0.5% by mass or more and 2% by mass or less, including Si of 0.5% by mass or less, Mn of 0.5% by mass or less, and N of 0.003% by mass or more and 0.04% by mass or less. The balance is preferably composed of Co and impurities. Impurities may include Al of 0.5% by mass or less and O of 0.04% by mass or less.

C: 0.08% by mass or more and 0.25% by mass or less
The C component is an MC-type carbide phase (a carbide phase of Ti, Zr, Hf, V, Nb and / or Ta, which may be hereinafter referred to as a precipitation-hardened carbide phase) to be a precipitation-hardened phase and a grain boundary of a parent phase crystal. _{It is an important component that constitutes the M 23} C ₆ type carbide phase (Cr, Fe, W, Mo and / or Mn carbide phase, hereinafter sometimes referred to as grain boundary strengthened carbide phase), which is the phase that suppresses slippage. be. The content of the C component is preferably 0.08% by mass or more and 0.25% by mass or less, more preferably 0.1% by mass or more and 0.2% by mass or less, and further preferably 0.12% by mass or more and 0.18% by mass or less. When the C content is less than 0.08% by mass, the precipitation amount of the precipitation strengthened carbide phase and the grain boundary strengthening carbide phase is insufficient, and the effect of improving the mechanical properties cannot be sufficiently obtained. On the other hand, when the C content exceeds 0.25% by mass, the Cr carbide phase is excessively precipitated or excessively hardened, so that the ductility and toughness of the alloy material are lowered.

B: 0.1% by mass or less
The B component is a component that contributes to the improvement of the bondability of the crystal grain boundaries (so-called grain boundary strengthening). The component B is not an essential component, but when it is contained, it is preferably 0.1% by mass or less, more preferably 0.005% by mass or more and 0.05% by mass or less. When the B content exceeds 0.1% by mass, cracks (for example, solidification cracks) are likely to occur during the formation of the AM body.

Cr: 10% by mass or more and 30% by mass or less
The Cr component is a component that contributes to the improvement of corrosion resistance and oxidation resistance, and is a component that constitutes a grain boundary reinforced carbide phase. The content of the Cr component is preferably 10% by mass or more and 30% by mass or less, and more preferably 15% by mass or more and 27% by mass or less. When a corrosion-resistant coating layer is separately provided on the outermost surface of the Co-based alloy product, the content of the Cr component is more preferably 10% by mass or more and 18% by mass or less. When the Cr content is less than 10% by mass, the action effect (improvement of corrosion resistance and oxidation resistance, formation of grain boundary strengthened carbide phase) cannot be sufficiently obtained. On the other hand, when the Cr content exceeds 30% by mass, the brittle σ phase is formed or the Cr carbide phase is excessively formed, and the mechanical properties (toughness, ductility, strength) are lowered.

Ni: 30% by mass or less
Since the Ni component has properties similar to those of the Co component and is cheaper than the Co component, it is a component that can be contained in a form that replaces a part of the Co component. The Ni component is not an essential component, but when it is contained, it is preferably 30% by mass or less, more preferably 20% by mass or less, and further preferably 5% by mass or more and 15% by mass or less. When the Ni content exceeds 30% by mass, the wear resistance and resistance to local stress, which are the characteristics of Co-based alloys, decrease. This is considered to be due to the difference between the stacking defect energy of Co and that of Ni.

Fe: 5% by mass or less
Since the Fe component is much cheaper than the Ni component and has properties similar to those of the Ni component, it is a component that can be contained in a form that replaces a part of the Ni component. It is also a component that can constitute a grain boundary-enhanced carbide phase. The total content of Fe and Ni is preferably 30% by mass or less, more preferably 20% by mass or less, and further preferably 5% by mass or more and 15% by mass or less. The Fe component is not an essential component, but when it is contained, it is preferably 5% by mass or less, more preferably 3% by mass or less in a range smaller than the Ni content. When the Fe content exceeds 5% by mass, it becomes a factor of deterioration of corrosion resistance and mechanical properties.

W and / or Mo: Total 5% by mass or more and 12% by mass or less
The W component and the Mo component are components that contribute to the solid solution strengthening of the matrix phase, and some of them are also components that can constitute the grain boundary strengthening carbide phase. The total content of the W component and / or the Mo component (one or more of the W component and the Mo component) is preferably 5% by mass or more and 12% by mass or less, and more preferably 7% by mass or more and 10% by mass or less. When the total content of the W component and the Mo component is less than 5% by mass, the solid solution strengthening of the matrix phase becomes insufficient. On the other hand, when the total content of the W component and the Mo component exceeds 12% by mass, the brittle σ phase is likely to be generated and the mechanical properties (toughness, ductility) are lowered.

Re: 2% by mass or less
The Re component is a component that contributes to the strengthening of the solid solution of the matrix and the improvement of corrosion resistance. The Re component is not an essential component, but when it is contained, it is preferably 2% by mass or less in the form of replacing a part of the W component or the Mo component, and more preferably 0.5% by mass or more and 1.5% by mass or less. When the Re content exceeds 2% by mass, the action and effect of the Re component is saturated, and the increase in material cost becomes a demerit.

One or more of Ti, Zr, Hf, V, Nb and Ta: 0.5% by mass or more and 2% by mass or less in total
The Ti component, Zr component, Hf component, V component, Nb component and Ta component are important components constituting the precipitation-hardened carbide phase (MC type carbide phase). The total content of one or more of the Ti, Zr, Hf, V, Nb and Ta components is preferably 0.5% by mass or more and 2% by mass or less, and more preferably 0.5% by mass or more and 1.8% by mass or less. If the total content is less than 0.5% by mass, the amount of precipitation-hardened carbide phase precipitated is insufficient, and the effect of improving mechanical properties cannot be sufficiently obtained. On the other hand, when the total content exceeds 2% by mass, the precipitation-hardened carbide phase particles become coarse, promote the formation of a brittle phase (for example, σ phase), or generate oxide phase particles that do not contribute to precipitation strengthening. The mechanical properties are reduced.

More specifically, when Ti is contained, the content is preferably 0.01% by mass or more and 1% by mass or less, and more preferably 0.05% by mass or more and 0.8% by mass or less.

When Zr is contained, the content is preferably 0.05% by mass or more and 1.5% by mass or less, and more preferably 0.1% by mass or more and 1.2% by mass or less. From the viewpoint of mechanical strength, it is more preferable to use the Zr component as an essential component.

When Hf is contained, the content is preferably 0.01% by mass or more and 0.5% by mass or less, and more preferably 0.02% by mass or more and 0.1% by mass or less.

When V is contained, the content is preferably 0.01% by mass or more and 0.5% by mass or less, and more preferably 0.02% by mass or more and 0.1% by mass or less.

When Nb is contained, the content is preferably 0.02% by mass or more and 1% by mass or less, and more preferably 0.05% by mass or more and 0.8% by mass or less.

When Ta is contained, the content is preferably 0.05% by mass or more and 1.5% by mass or less, and more preferably 0.1% by mass or more and 1.2% by mass or less.

Si: 0.5% by mass or less
The Si component is a component that plays a role of deoxidizing and contributes to the improvement of mechanical properties. The Si component is not an essential component, but when it is contained, it is preferably 0.5% by mass or less, and more preferably 0.01% by mass or more and 0.3% by mass or less. When the Si content exceeds 0.5% by mass _{, coarse particles of oxide (for example, SiO 2} ) are formed, which causes deterioration of mechanical properties.

Mn: 0.5% by mass or less
The Mn component is a component that plays a role of deoxidation and desulfurization and contributes to improvement of mechanical properties and corrosion resistance. It is also a component that can constitute a grain boundary-enhanced carbide phase. The Mn component is not an essential component, but when it is contained, it is preferably 0.5% by mass or less, and more preferably 0.01% by mass or more and 0.3% by mass or less. When the Mn content exceeds 0.5% by mass, coarse particles of sulfide (for example, MnS) are formed, which causes deterioration of mechanical properties and corrosion resistance.

N: 0.003% by mass or more and 0.04% by mass or less
The N component is a component that contributes to the stable formation of the precipitation-hardened carbide phase. The content of the N component is preferably 0.003% by mass or more and 0.04% by mass or less, more preferably 0.005% by mass or more and 0.03% by mass or less, and further preferably 0.007% by mass or more and 0.025% by mass or less. If the N content is less than 0.003% by mass, the action and effect of the N component cannot be sufficiently obtained. On the other hand, when the N content exceeds 0.04% by mass, coarse particles of nitride (for example, Cr nitride) are formed, which causes a decrease in mechanical properties.

Remaining: Co component + impurities
The Co component is one of the main components of this alloy and is the component having the maximum content. As described above, the Co-based alloy material has an advantage of having corrosion resistance and wear resistance equal to or higher than that of the Ni-based alloy material.

The Al component is one of the impurities of this alloy and is not a component intentionally contained. However, if the Al content is 0.5% by mass or less, it is acceptable because it does not significantly adversely affect the mechanical properties of the Co-based alloy product. When the Al content exceeds 0.5% by mass _{, coarse particles of oxides and nitrides (for example, Al 2} O ₃ and Al N) are formed, which causes deterioration of mechanical properties.

The O component is also one of the impurities of this alloy and is not a component intentionally contained. However, an O content of 0.04% by mass or less is acceptable because it does not significantly adversely affect the mechanical properties of the Co-based alloy product. When the O content exceeds 0.04% by mass, coarse particles of various oxides (for example, Ti oxide, Zr oxide, Al oxide, Fe oxide, Si oxide) are formed, which causes deterioration of mechanical properties. become.

In this step S1, there is no particular limitation on the method / method for preparing the Co-based alloy powder, and the conventional method / method can be used. For example, a mother alloy ingot production element step (S1a) for producing a mother alloy ingot (master ingot) by mixing, melting, and casting raw materials so as to have a desired chemical composition, and forming an alloy powder from the mother alloy ingot. The atomizing elementary process (S1b) may be performed. In addition, there are no particular restrictions on the atomizing method, and conventional methods and methods can be used. For example, a gas atomizing method or a centrifugal atomizing method that can obtain high-purity, spherical particles can be preferably used.

The particle size of the alloy powder is preferably 5 μm or more and 100 μm or less, more preferably 10 μm or more and 70 μm or less, and 10 μm or more and 50 μm or less from the viewpoint of handleability and filling property of the alloy powder bed in the selective laser melting step S2 of the next step. More preferred. If the particle size of the alloy powder is less than 5 μm, the fluidity of the alloy powder decreases in the next step S2 (the formability of the alloy powder bed decreases), which causes the shape accuracy of the AM body to decrease. On the other hand, when the particle size of the alloy powder exceeds 100 μm, it becomes difficult to control the local melting and rapid solidification of the alloy powder bed in the next step S2, the melting of the alloy powder becomes insufficient, and the surface roughness of the AM body increases. It becomes a factor to do.

From the above, it is preferable to carry out the alloy powder classifying step (S1c) for classifying the particle size of the alloy powder into the range of 5 μm or more and 100 μm or less. In the present invention, even if it is confirmed that the particle size distribution of the alloy powder produced in the atomizing elementary step S1b is within a desired range as a result of measuring the particle size distribution, it is considered that the present elementary step S1c has been performed.

(Selective laser melting process)
This step S2 is a step of forming an AM body having a desired shape by a selective laser melting (SLM) method using the prepared Co-based alloy powder. Specifically, an alloy powder bed preparation step (S2a) in which a Co-based alloy powder is spread to prepare an alloy powder bed having a predetermined thickness, and a predetermined region of the alloy powder bed is irradiated with laser light to prepare the alloy powder bed in the region. This is a process of forming an AM body by repeating the laser melt coagulation element step (S2b) in which the Co-based alloy powder is locally melted and rapidly solidified.

In this step S2, in order to obtain a desired microstructure (fine structure in which precipitation-hardened carbide phase particles are dispersed and precipitated in the parent phase crystal grains) in the final Co-based alloy product, it becomes a precursor of the product. Controls the microstructure of the AM body. Then, in order to control the microstructure of the AM body, local melting and rapid solidification of the alloy powder bed are controlled.

More specifically, in relation to the thickness h (unit: μm) of the alloy powder bed, the output P (unit: W) of the laser beam, and the scanning speed S (unit: mm / s) of the laser beam, “15 The thickness h of the alloy powder bed, the laser light output P, and the laser light scanning speed S so as to satisfy <h <150 "and" 67 (P / S) -3.5 <h <2222 (P / S) +13 ". It is preferable to control. If the control conditions are not met, an AM body having a desired microstructure cannot be obtained.

The output P of the laser beam and the scanning speed S of the laser beam basically depend on the configuration of the laser device, but should be selected within the range of "10 ≤ P ≤ 1000" and "10 ≤ S ≤ 7000", for example. Just do it.

(Co-based alloy laminated model)
FIG. 2 is a scanning electron microscope (SEM) observation image showing an example of the microstructure of the Co-based alloy AM body obtained in the SLM step S2. As shown in FIG. 2, the Co-based alloy AM body obtained in the SLM step S2 has an extremely specific microstructure that has never been seen before.

The AM body is a polycrystal of a parent phase crystal, and segregated cells having an average size of 0.13 μm or more and 2 μm or less are formed in the crystal grains of the polycrystal. The average size of the segregated cell is more preferably 0.15 μm or more and 1.5 μm or less from the viewpoint of mechanical strength. It is confirmed that carbide phase particles may be precipitated in a part of the boundary region of the segregation cell. In addition, from many experiments, it was confirmed that the average crystal grain size of the matrix crystal is preferably 5 μm or more and 150 μm or less.

In the present invention, the size of the segregated cell is basically defined as the average of the major axis and the minor axis, but when the aspect ratio between the major axis and the minor axis is 3 or more, twice the minor axis is adopted. It shall be. Further, the average spacing of the particles of the precipitation strengthened carbide phase in the present invention is defined as being represented by the size of the segregation cell because the particles are precipitated on the boundary region of the segregation cell.

A more detailed microstructure observation was performed using a scanning transmission electron microscope-energy dispersive X-ray spectroscopy (STEM-EDX). , It was confirmed that the components (Ti, Zr, Hf, V, Nb, Ta, C) forming the precipitation-enhanced carbide phase were segregated in the cell wall-like region. It was also confirmed that the particles precipitated on the boundary region of the segregation cell were MC-type carbide phase particles.

As a matter of course, this AM body also has segregation of components forming the MC-type carbide phase and precipitation of MC-type carbide phase particles on the grain boundaries of the parent phase crystal grains.

(First heat treatment step)
This step S3 is a step of performing the first heat treatment on the formed Co-based alloy AM body. As the condition of the first heat treatment, heat treatment in a temperature range of 750 ° C. or higher and lower than 1100 ° C. is preferable. The heat treatment temperature is more preferably 800 ° C. or higher and 1050 ° C. or lower, and further preferably 850 ° C. or higher and 1000 ° C. or lower. The holding time in the heat treatment may be appropriately set in the range of 0.5 hours or more and 10 hours or less in consideration of the temperature. The cooling method after the heat treatment is not particularly limited, and for example, oil cooling, water cooling, air cooling, or furnace cooling may be used.

FIG. 3 is an SEM observation image showing an example of the fine structure of the Co-based alloy laminated model body subjected to the first heat treatment. As shown in FIG. 3, the Co-based alloy AM body subjected to the first heat treatment also has an extremely specific microstructure that has never been seen before.

Very interestingly, by performing the first heat treatment, the components segregated in the boundary region of the segregation cell diffuse and combine on the boundary (along the boundary) to form a precipitation strengthened carbide phase (MC type carbide phase). It was found that the cell wall of the segregated cell almost disappeared when it started to form (more accurately, it becomes difficult to confirm the cell wall of the segregated cell by microstructure observation). In other words, the particles of the precipitation strengthened carbide phase are dispersed and formed along the region where the cell wall would have been (above the boundary region of the original segregation cell). In the present invention, a region surrounded by precipitation-hardened carbide phase particles precipitated along a region where the cell wall of such a segregation cell would have been located will be referred to as a “post-segregation cell”.

The shape of the post-segregation cell is considered to be almost the same as the shape of the segregation cell, and the average size of the post-segregation cell is 0.13 to 2 μm. Since the precipitated precipitation-strengthened carbide phase particles can serve as a pinning point for the grain boundary movement of the matrix crystal grains, the coarsening of the matrix crystal grains is suppressed.

It was also found that in addition to the formation of the precipitation-strengthened carbide phase, the grain boundary-reinforced carbide phase (M ₂₃ C ₆ type carbide phase) also began to form on the grain boundaries of the matrix crystal.

Furthermore, by performing the first heat treatment, it is possible to alleviate the residual internal strain of the AM body that may occur during the rapid solidification of the SLM step S2, and undesired deformation during the post-step or use of the alloy product. Can be prevented.

(Second heat treatment step)
This step S4 is a step of performing a second heat treatment on the Co-based alloy AM body that has been subjected to the first heat treatment. As the conditions for the second heat treatment, heat treatment at a temperature lower than that of the previous first heat treatment in a temperature range of 600 ° C. or higher and 1000 ° C. or lower is preferable. The holding time in the heat treatment may be appropriately set in the range of 0.5 hours or more and 20 hours or less in consideration of the temperature. The cooling method after the heat treatment is not particularly limited, and for example, oil cooling, water cooling, air cooling, or furnace cooling may be used.

FIG. 4 is an SEM observation image showing an example of the fine structure of the Co-based alloy product obtained in the second heat treatment step S4. As shown in FIG. 4, the Co-based alloy product obtained in the second heat treatment step S4 also has an extremely specific microstructure that has never been seen before. By performing the second heat treatment, the precipitation-strengthened carbide phase particles and the grain boundary-reinforced carbide phase particles can be appropriately grown while suppressing the coarsening of the matrix crystal grains.

As a result of STEM-EDX analysis, the precipitation strengthened carbide phase can be regarded as the MC type carbide phase containing Ti, Zr, Hf, V, Nb and / or Ta, and the grain boundary strengthened carbide phase is Cr, Fe, W. It was confirmed that it can be regarded _{as an M 23} C ₆ type carbide phase containing, Mo and / or Mn.

As a result of step S4, in the Co-based alloy product, the average crystal grain size of the matrix crystals was 5 μm or more and 150 μm or less, and precipitation-hardened carbides were precipitated in the grains of each matrix crystal at an average interparticle distance of 0.13 μm or more and 2 μm or less. It has a microstructure in which the phase particles are dispersed and precipitated, and the particles of the grain boundary reinforced carbide phase are precipitated on the grain boundaries of the matrix crystal. In the Co-based alloy product of the present invention, as described above, the particles of the precipitation-strengthened carbide phase are dispersed and precipitated on the grain boundaries of the parent phase crystals.

[Co-based alloy product]
FIG. 5 is an example of a Co-based alloy product according to the present invention, and is a schematic perspective view showing a turbine vane as a high temperature member. As shown in FIG. 5, the turbine stationary blade 100 is roughly composed of an inner ring side end wall 101, a blade portion 102, and an outer ring side end wall 103. A cooling structure is often formed inside the wing. As described above, since the turbine vane 100 has a very complicated shape and structure, the technical significance of the AM body formed by the near net shape and the alloy product based on the AM body is great.

For example, in the case of a gas turbine for power generation with an output of 30 MW class, the length of the blade of the turbine stationary blade (distance between both end walls) is about 170 mm. Further, the Co-based alloy product of the present invention may, of course, be used as a turbine blade.

FIG. 6 is a schematic cross-sectional view showing an example of a gas turbine equipped with the Co-based alloy product according to the present invention. As shown in FIG. 6, the gas turbine 200 is roughly composed of a compressor unit 210 that compresses intake air and a turbine unit 220 that blows fuel combustion gas onto turbine blades to obtain rotational power. The high temperature member of the present invention can be suitably used as a turbine nozzle 221 or a turbine stationary blade 100 in the turbine section 220. Further, the high temperature member of the present invention is not limited to gas turbine applications, but may be used for other turbine applications (for example, steam turbine applications), or is used in a high temperature environment in other machines / devices. It may be a member.

FIG. 7 is an example of a Co-based alloy product according to the present invention, and is a schematic perspective view showing a heat exchanger as a high temperature member. The heat exchanger 300 shown in FIG. 7 is an example of a plate fin type heat exchanger, and basically has a structure in which separate layers 301 and fin layers 302 are alternately laminated. Both ends of the fin layer 302 in the flow path width direction are sealed by the sidebar portion 303. By alternately circulating the high temperature fluid and the low temperature fluid through the adjacent fin layer 302, heat exchange is performed between the high temperature fluid and the low temperature fluid.

The heat exchanger 300 according to the present invention is integrally formed without brazing or welding the components (for example, separate plates, corrugated fins, sidebars) in the conventional heat exchanger. It can be made more heat resistant and lighter than a heat exchanger. Further, by forming an appropriate uneven shape on the surface of the flow path, the fluid can be turbulent and the heat transfer efficiency can be improved. Improving heat transfer efficiency leads to miniaturization of heat exchangers.

Hereinafter, the present invention will be described in more detail with reference to Examples and Comparative Examples. The present invention is not limited to these examples.

[Experiment 1]
(Preparation of alloy powders IA-1 to IA-7 and CA-1 to CA-5)
A Co-based alloy powder having the chemical composition shown in Table 1 was prepared (alloy powder preparation step S1). Specifically, first, the raw materials were mixed, and then melted and cast by a vacuum high-frequency induction melting method to prepare a mother alloy ingot (mass: about 2 kg). Next, the atomizing element step S1b was performed in which the mother alloy ingot was redissolved to form an alloy powder by a gas atomizing method in an argon gas atmosphere.

Next, for each of the obtained alloy powders, an alloy powder classification element step S1c for controlling the particle size of the alloy powder was performed to classify the powder particle size in the range of 5 to 25 μm.

As shown in Table 1, the invention alloy powders IA-1 to IA-7 are alloy powders having a chemical composition satisfying the provisions of the present invention. On the other hand, the comparative alloy powder CA-1 has a C content and a Cr content outside the provisions of the present invention. In the comparative alloy powder CA-2, the C content, the Ni content, and the total content of "Ti + Zr + Hf + V + Nb + Ta" are out of the provisions of the present invention. In the comparative alloy powder CA-3, the C content, the N content, and the total content of "Ti + Zr + Hf + V + Nb + Ta" are out of the provisions of the present invention. The total content of "Ti + Zr + Hf + V + Nb + Ta" of the comparative alloy powder CA-4 is out of the specification of the present invention. In the comparative alloy powder CA-5, the W content and the total content of "Ti + Zr + Hf + V + Nb + Ta" are out of the scope of the present invention.

[Experiment 2]
(Examination of SLM conditions in selective laser melting process)
An AM body (diameter 8 mm × length 10 mm) was formed by the SLM method using the IA-4 alloy powder prepared in Experiment 1 (selective laser melting step S2). The SLM condition is that the output P of the laser beam is 85 W, and the local heat input P / S (unit: W. s / mm = J / mm) was controlled. Controlling the amount of local heat input corresponds to controlling the cooling rate.

The average size of the segregated cells was measured by observing the microstructure of each AM body prepared above. Microstructure observation was performed by SEM. In addition, the average size of segregated cells was measured by image analysis of the obtained SEM observation image using image processing software (ImageJ, public domain software developed by the National Institutes of Health (NIH) in the United States).

FIG. 8 is an example of SLM conditions in the selective laser melting step S2, and is a graph showing the relationship between the thickness of the alloy powder bed and the amount of local heat input. In FIG. 8, as a result of observing the microstructure of the formed AM body, those having an average size of segregated cells in the range of 0.15 to 1.5 μm were judged as “passed” and indicated by “◯” in the figure, and other than that. Those were judged to be "failed" and indicated by "x" in the figure.

From the results of Experiment 2, the SLM conditions in the selective laser melting step S2 are the thickness h (unit: μm) of the alloy powder bed, the laser light output P (unit: W), and the laser light scanning speed S (unit: unit:). It is confirmed that it is preferable to control the relationship with mm / s) so as to satisfy "15 <h <150" and "67 (P / S) -3.5 <h <2222 (P / S) +13". .. That is, the hatched area is the pass determination area.

[Experiment 3]
(Examination of heat treatment conditions in the first heat treatment step)
Using the alloy powders of IA-5 and IA-6 prepared in Experiment 1, an AM body (diameter 10 mm × length 50 mm) was formed by the SLM method (selective laser melting step S2). The SLM conditions are that the thickness h of the alloy powder bed is 100 μm, the output P of the laser beam is 100 W, and the local heat input P / S (unit: unit:) by controlling the scanning speed S (mm / s) of the laser beam. W · s / mm = J / mm) was controlled and adjusted to satisfy the passing conditions of Experiment 2.

Each AM body prepared above was heat-treated at 950 to 1250 ° C. for 2 to 4 hours (first heat treatment step S3). Next, each AM body subjected to the first heat treatment is subjected to a heat treatment held at 900 ° C. for 4 hours (second heat treatment step S4), and a Co-based alloy using the IA-5 powder and the IA-6 powder is used. Products (IAP-5a to IAP-5d, IAP-6a to IAP-6d) were prepared. Test pieces for mechanical property test were collected from each of the produced alloy products and subjected to mechanical property test.

As a mechanical property test, a creep test was performed under the conditions of a temperature of 850 ° C. and a stress of 168 MPa, and the creep rupture time was measured. From the required characteristics for the high temperature member targeted by the present invention, a creep rupture time of 1200 hours or more was determined as "pass", and a creep rupture time of less than 1200 hours was determined as "fail". This creep characteristic can be said to be the same creep characteristic as the Ni-based alloy material.

The results of Experiment 3 are shown in Table 2.

As shown in Table 2, the samples under the condition that the first heat treatment temperature was 1100 ° C. or higher did not pass the mechanical properties. On the other hand, all the samples under the condition that the first heat treatment temperature was less than 1100 ° C. passed the mechanical properties. This is because the average particle size of the parent phase crystal grains is in an appropriate range, and the precipitation strengthened carbide phase particles (MC type carbide phase particles) and the grain boundary strengthened carbon dioxide phase particles (M ₂₃ C ₆ type carbide phase particles). ) Is considered to be due to the dispersion precipitation in a preferable form.

[Experiment 4]
(Alloy products using IA-1 to IA-7 powder IAP-1-1 to IAP-7-1, and alloy products using CA-1 to CA-5 powder CAP-1-1 to CAP- Production of 5-1)
Using the alloy powders of IA-1 to IA-7 and CA-1 to CA-5 prepared in Experiment 1, an AM body (diameter 10 mm x length 50 mm) was formed by the SLM method in the same manner as in Experiment 3. (Selective laser melting step S2).

Each AM body produced above was heat-treated at 1000 ° C. for 4 hours (first heat treatment step S3). Next, each AM body subjected to the first heat treatment is heat-treated by holding it at 850 ° C. for 1 hour (second heat treatment step S4) to produce a Co-based alloy using IA-1 to IA-7 powders. Co-based alloy products CAP-1-1 to CAP-5-1 using the products IAP-1-1 to IAP-7-1 and CA-1 to CA-5 powders were prepared.

(Microstructure observation and mechanical property test)
From the Co-based alloy products IAP-1-1 to IAP-7-1 and CAP-1-1 to CAP-5-1 prepared above, test pieces for microstructure observation and mechanical property test were collected, respectively. Then, microstructure observation and mechanical property test were performed.

For microstructure observation, the same SEM observation and image analysis of the SEM observation image as in Experiment 2 were performed to determine the presence or absence of co-precipitation of the precipitation-hardened carbide phase particles and the grain boundary-hardened carbide phase particles, and the precipitation strengthening in the parent phase crystal grains. The average intergranular distance of the carbide phase particles was investigated. Further, as the mechanical property test, the same creep test as in Experiment 3 was performed, and "pass / fail" was judged by the same criteria as in Experiment 3. The results of Experiment 4 are shown in Table 3.

As shown in Table 3, in the alloy products IAP-1-1 to IAP-7-1 according to the present invention, precipitation-hardened carbide phase particles and grain boundary-hardened carbide phase particles are co-precipitated, and the parent phase crystal grains. It was confirmed that the average intergranular distance of the precipitation-hardened carbide phase particles was in the range of 0.15 to 1.5 μm. In addition, it was confirmed that all of these samples passed the mechanical properties.

On the other hand, in the alloy products CAP-1-1 to CAP-3-1, precipitation strengthened carbide phase particles and grain boundary strengthened carbide phase particles were co-precipitated, but the creep test failed in all the samples. .. When viewed individually, it is considered that CAP-1-1 is caused by excessive precipitation of Cr carbide particles because the C content and the Cr content are excessive. Since the C content and the total content of "Ti + Zr + Hf + V + Nb + Ta" are excessive in CAP-2-1, it is considered that the average interparticle distance is increased as a result of the coarsening of the precipitation-hardened carbide phase particles. Be done. In CAP-3-1, the C content was excessive and the N content and the total content of "Ti + Zr + Hf + V + Nb + Ta" were too small. It is considered that this is due to the excessive distance between particles. From the results of CAP-1-1 to CAP-3-1, if Cr carbide particles are excessively precipitated or the average interparticle distance of the precipitation strengthened carbide phase is excessive, the mechanical properties cannot secure the required level. It is confirmed.

In addition, since the total content of "Ti + Zr + Hf + V + Nb + Ta" in CAP-4-1 and CAP-5-1 is too small (because it is hardly contained), the stage of AM after the selective laser melting step S2. The segregation cell was not formed in the above process, and the precipitation-strengthened carbide phase particles were not dispersed and precipitated in the parent phase crystal grains even after the first heat treatment step S3 and the second heat treatment step S4, resulting in poor mechanical properties. It is probable that it passed.

From the results of Experiment 4, it was confirmed that IA-1 to IA-7 having the chemical composition specified in the present invention are preferable as the starting powder of the Co-based alloy product. Further, Co. It was confirmed that the creep characteristics of the base alloy product could be improved.

The above-described embodiments and experimental examples have been described for the purpose of assisting the understanding of the present invention, and the present invention is not limited to the specific configurations described. For example, it is possible to replace a part of the configuration of the embodiment with the configuration of the common general technical knowledge of those skilled in the art, and it is also possible to add the configuration of the common general technical knowledge of the person skilled in the art to the configuration of the embodiment. That is, the present invention may delete, replace with another configuration, or add another configuration to a part of the configurations of the embodiments and experimental examples of the present specification without departing from the technical idea of the invention. It is possible.

100 ... Turbine stationary blade, 101 ... Inner ring side end wall, 102 ... Blade part, 103 ... Outer ring side end wall, 200 ... Gas turbine, 210 ... Compressor part, 220 ... Turbine part, 221 ... Turbine nozzle, 300 ... Heat exchange Vessel, 301 ... Separate layer, 302 ... Fin layer, 303 ... Sidebar part.

Claims (7)

A product made of a cobalt-based alloy
The cobalt-based alloy is
With carbon of 0.08% by mass or more and 0.25% by mass or less,
Boron of 0.1% by mass or less and
Contains 10% by mass or more and 30% by mass or less of chromium
It contains 5% by mass or less of iron and 30% by mass or less of nickel, and the total of the iron and the nickel is 30% by mass or less.
It contains tungsten and / or molybdenum, and the total of the tungsten and the molybdenum is 5% by mass or more and 12% by mass or less.
Titanium, and zirconium, hafnium, vanadium, comprises less than 2 wt% of two or more total 0.5 wt% or more of niobium and tantalum, as essential following zirconium 1.5 wt% 0.05 wt%,
With 0.5% by mass or less of silicon,
With 0.5% by mass or less of manganese,
Contains 0.003% by mass or more and 0.04% by mass or less of nitrogen.
The rest consists of cobalt and impurities,
The impurities are
With 0.5% by mass or less of aluminum,
Has a chemical composition containing less than 0.04% by weight oxygen
The product is a polycrystal of parent phase crystal grains having an average crystal grain size of 5 μm or more and 150 μm or less , and MC-type carbide phase particles and M ₂₃ C _6- type carbide phase particles are co-precipitated.
The particles of the MC-type carbide phase are dispersed and precipitated in the grains of the parent phase crystal grains at an average interparticle distance of 0.13 μm or more and 2 μm or less.
The particles of the M ₂₃ C ₆ type carbide phase are precipitated on the grain boundaries of the parent phase crystal grains.
A cobalt-based alloy product characterized by this. In the cobalt-based alloy product according to claim 1,
The particles of the MC-type carbide phase are particles of the MC-type carbide phase containing the titanium, the zirconium, the hafnium, the vanadium, the niobium and / or the tantalum.
Particles of the M ₂₃ C ₆ type carbide phase, the chromium, the iron, cobalt based alloys produced, wherein the tungsten particles of the molybdenum and / or M ₂₃ C ₆ type carbide phase comprising said manganese thing. In the cobalt-based alloy product according to claim 1 or 2.
The chemical composition of the cobalt-based alloy is
When the titanium is contained, the titanium is 0.01% by mass or more and 1% by mass or less .
If the previous SL containing hafnium, said hafnium is 0.5 mass% 0.01 mass% or more,
When the vanadium is contained, the vanadium is 0.01% by mass or more and 0.5% by mass or less.
When the niobium is contained, the niobium is 0.02% by mass or more and 1% by mass or less.
When the tantalum is contained, the tantalum is 0.05% by mass or more and 1.5% by mass or less.
A cobalt-based alloy product characterized by this. In the cobalt-based alloy product according to any one of claims 1 to 3.
The product is a cobalt-based alloy product characterized by having a creep rupture time of 1200 hours or more when a creep test is performed under the conditions of a temperature of 850 ° C. and a stress of 168 MPa. In the cobalt-based alloy product according to any one of claims 1 to 4.
The product is a cobalt-based alloy product characterized by being a high-temperature member. In the cobalt-based alloy product according to claim 5.
A cobalt-based alloy product, wherein the high-temperature member is a turbine blade, a turbine blade, a turbine combustor nozzle, or a heat exchanger. The method for producing a cobalt-based alloy product according to any one of claims 1 to 6.
An alloy powder preparation step for preparing a cobalt-based alloy powder having the chemical composition, and
An alloy powder bed preparation step of spreading the cobalt-based alloy powder to prepare an alloy powder bed having a predetermined thickness, and irradiating a predetermined region of the alloy powder bed with laser light to obtain the cobalt-based alloy powder in the region. A laser melting and coagulating element step of locally melting and rapid solidifying, and a selective laser melting step of forming a laminated model by repeating the steps of
The first heat treatment step of performing the first heat treatment in the temperature range of 750 ° C. or higher and lower than 1100 ° C. on the laminated model, and the first heat treatment step.
It has a second heat treatment step of subjecting the laminated model body subjected to the first heat treatment to a second heat treatment having a temperature lower than that of the first heat treatment in a temperature range of 600 ° C. or higher and 1000 ° C. or lower.
The alloy powder preparation step includes an alloy powder classifying element step for classifying the cobalt-based alloy powder into a particle size range of 5 μm or more and 100 μm or less.
In the selective laser melting step, the predetermined thickness h (unit: μm) of the alloy powder bed, the output P (unit: W) of the laser beam, and the scanning speed S (unit: mm / s) of the laser beam. The predetermined thickness h, the output P, and the scanning speed so that the relationship with the above satisfies “15 <h <150” and “67 (P / S) −3.5 <h <2222 (P / S) +13”. Control with S,
A method for producing a cobalt-based alloy product.

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